High-strength steel sheet having excellent impact resistant property and method for manufacturing thereof

ABSTRACT

Provided is a method for manufacturing a steel sheet, the method including: reheating a steel slab at a temperature of 1200° C. to 1350° C., the steel slab including, by weight%, carbon (C): 0.05% to 0.14%, silicon (Si): 0.01% to 1.0%, manganese (Mn): 1.5% to 2.5%, aluminum (Al): 0.01% to 0.1%, chromium (Cr): 0.005% to 1.0%, phosphorus (P): 0.001% to 0.05%, sulfur (S): 0.001% to 0.01%, nitrogen (N): 0.001% to 0.01%, niobium (Nb): 0.005% to 0.06%, titanium (Ti): 0.005% to 0.11%, and a balance of iron (Fe) and inevitable impurities; finish hot rolling the reheated steel slab under predetermined conditions to obtain a hot-rolled steel sheet; cooling the hot-rolled steel sheet at a cooling rate of 10° C./s to 100° C./s to a temperature of 400° C. to 500° C. after the finish hot rolling; and coiling the steel sheet at a temperature of 400° C. to 500° C. after the cooling.

TECHNICAL FIELD

The present disclosure relates to a steel sheet for automobile chassisparts, and more particularly, to a high-strength steel sheet having highimpact resistance property and a method for manufacturing thehigh-strength steel sheet.

BACKGROUND ART

High-strength hot-rolled steel sheets are mainly used for automotiveparts such as chassis members, lower arms, reinforcing materials, orconnecting materials, and techniques for improving the formability ofsuch high-strength hot-rolled steel sheets have been proposed.

For example, techniques for improving the stretch-flangeability of steelsheets by forming the steel sheets as complex-phase steel sheets havinga ferrite-bainite dual phase as a microstructure, or techniques formanufacturing steel sheets having high-strength and high-burringproperties by forming ferrite or bainite as a matrix have been proposed.

Specifically, in Patent Document 1, a steel sheet is maintained in aferrite transformation zone for several seconds under specific coolingconditions immediately after hot rolling, and then the steel sheet iscoiled at a bainite formation temperature such that bainite is formed inthe steel sheet, thereby forming a complex phase of polygonal ferriteand bainite as the microstructure of the steel sheet and guaranteeingthe strength and stretch-flangeability of the steel sheet.

In addition, Patent Document 2 discloses a high-burring steel havingC-Si-Mn-based bainitic ferrite and granular bainitic ferrite as amatrix, and Patent Document 3 discloses a method for improving thestretch-flangeability of a steel sheet by forming bainite in an areafraction of 95% or more and minimizing grains stretched in the rollingdirection.

To manufacture the above-mentioned high-strength steels, alloyingelements such as Si, Mn, Al, Mo, and Cr are mainly used, which areeffective in improving the strength and stretch-flangeability ofhot-rolled steel sheets.

However, it is necessary to add such alloying elements in large amountsto improve the above-mentioned physical properties, causing segregationof the alloying elements, non-uniform microstructures, or the like andthus results in deterioration in stretch-flangeability. In addition,automobiles using steels containing large amounts of such alloyingelements have poor crashworthiness because when automobiles crash,fractures easily occur in regions in which the alloying elements aresegregated and regions having non-uniform microstructures.

In particular, the microstructures of high-hardenability steelssensitively change during cooling, causing the formation of unevenlow-temperature transformation phases which results in deterioration inimpact resistance.

In addition, if precipitate-forming elements such as Ti, Nb, and V areexcessively added to increase the strength of steel, although the impactresistance of the steel increases owing to the increase in strength, theformability of the steel deteriorates because recrystallization isgreatly delayed during hot rolling and thus microstructures arestretched in the rolling direction.

-   (Patent Document 1) Japanese Patent Laid-Open Publication No.    1994-293910-   (Patent Document 2) Korean Patent Publication No. 10-1114672-   (Patent Document 3) Korean Patent Publication No. 10-1528084

DISCLOSURE Technical Problem

Aspects of the present disclosure are to provide a steel sheet havinghigh strength and high excellent impact resistance and a method formanufacturing the steel sheet.

The scope of the present disclosure is not limited to theabove-mentioned aspects. Other aspects of the present disclosure arestated in the following description, and the aspects of the presentdisclosure will be clearly understood by those of ordinary skill in theart through the following description.

Technical Solution

According to an aspect of the present disclosure, there is provided ahigh-strength steel sheet having high impact resistance, thehigh-strength steel sheet including, by weight%, carbon (C): 0.05% to0.14%, silicon (Si): 0.01% to 1.0%, manganese (Mn): 1.5% to 2.5%,aluminum (Al): 0.01% to 0.1%, chromium (Cr) : 0.005% to 1.0%, phosphorus(P) : 0.001% to 0.05%, sulfur (S): 0.001% to 0.01%, nitrogen (N): 0.001%to 0.01%, niobium (Nb) : 0.005% to 0.06%, titanium (Ti) : 0.005% to0.11%, and the balance of iron (Fe) and inevitable impurities,

wherein the high-strength steel sheet has a microstructure includingferrite and bainite in a total area fraction of 90% or more, and

the high-strength steel sheet has a value of 0.05 to 1.0 as a sheartexture ({110}<112>, {112)<111>) area ratio of a center region (rangingdeeper than 1/10t to ½t in a thickness direction, t refers to thickness(mm)) and a surface region (ranging from a surface to 1/10t in thethickness direction).

According to another aspect of the present disclosure, there is provideda method for manufacturing a high-strength steel sheet having highimpact resistance, the method including: reheating a steel slab havingthe above-described alloy composition at a temperature of 1200° C. to1350° C.; finish hot rolling the reheated steel slab under conditionssatisfying [Equation 1] and [Equation 2] below to obtain a hot-rolledsteel sheet; cooling the hot-rolled steel sheet at a cooling rate of 10°C./s to 100° C./s to a temperature of 400° C. to 500° C. after thefinish hot rolling; and coiling the steel sheet at a temperature of 400°C. to 500° C. after the cooling,

Tn-50≤FDT≤Tn   [Equation 1]

in which Tn refers to a temperature at which recrystallization delaystarts, Tn=730+92×[C]+70×[Mn]+45×[Cr]+780×[Nb]+520×[Ti]−80×[Si] whereeach element refers to a weight content, and FDT refers to a temperature(° C.) of the hot-rolled steel sheet immediately after the finish hotrolling,

Ec≤E1+E2≤1.5×Ec   [Equation 2]

in which Ec=4.75×10⁻⁴ (125×Exp(Qdef/(R×(273+FDT))))^(0.17),Qdef=277000−2535×[C]+1510×[Mn]+9621×[Si]+1255×[Cr]+53680×[Ti]^(0.592)+70730×[Nb]^(0.565)where each element refers to a weight content, and

E1 refers to a rolling reduction in a final pass of hot rolling, E2refers to a rolling reduction before the final hot rolling pass (passjust before the final pass), and R refers to the gas constant: 8.314.

Advantageous Effects

According to the present disclosure, it is possible to provide a steelsheet having high strength and high impact resistance. The steel sheetof the present disclosure may be suitably used as a material forautomobile chassis parts.

DESCRIPTION OF DRAWINGS

FIG. 1 is images of the shear texture of a surface region (A) and theshear texture of a center region (B) of Inventive Steel 6 according toan example of the present disclosure.

BEST MODE

The inventors have intensively studied changes in the strength andimpact properties of steel sheets according to the characteristics ofvarious alloying elements and microstructures of steels. As a result,the inventors have found that a steel sheet having high impactresistance and high strength can be obtained by appropriately adjustingthe contents of alloying elements in the steel sheet and optimizing thearea fractions of the matrix and shear texture of the steel sheet, andbased on this knowledge, the inventors have invented the presentinvention.

Hereinafter, the present disclosure will be described in detail.

According to an aspect of the present disclosure, a high-strength steelsheet having high impact resistance may preferably include, by weight %,carbon (C): 0.05% to 0.14%, silicon (Si): 0.01% to 1.0%, manganese (Mn):1.5% to 2.5%, aluminum (Al): 0.01% to 0.1%, chromium (Cr): 0.005% to1.0%, phosphorus (P): 0.001% to 0.05%, sulfur (S): 0.001% to 0.01%,nitrogen (N) : 0.001% to 0.01%, niobium (Nb) : 0.005% to 0.06%, titanium(Ti): 0.005% to 0.11%.

Hereinafter, reasons for limiting the alloy composition of the steelsheet will be described in detail. In this case, unless otherwisespecified, the content of each element is given in weight %.

Carbon (C): 0.05% to 0.14%

Carbon (C) is the most economical and effective element forstrengthening steel. As the content of C increases, precipitationstrengthening increases, and the fraction of bainite increases, therebyimproving tensile strength.

If the content of C in steel is lower than 0.05%, the steel may not besufficiently reinforced, whereas if the content of C exceeds 0.14%,martensite is formed to cause an excessive increase in strength anddeterioration in formability and impact resistance. In addition,weldability may also deteriorate.

Therefore, in the present disclosure, it is preferable that the contentof C be within the range of 0.05% to 0.14%. More preferably, the contentof C may be within the range of 0.06% to 0.13%.

Silicon (Si): 0.01% to 1.0%

Silicon (Si) has a function of deoxidizing molten steel and improvingstrength through solid-solution strengthening. In addition, Si isadvantageous in delaying the formation of coarse carbides for improvingthe formability of the steel sheet.

If the content of Si is lower than 0.01%, the effect of delaying theformation of carbides is insufficient, making it difficult to improveformability. Conversely, if the content of Si exceeds 1.0%, red scale isformed on the surface of the steel sheet during hot rolling, causing asignificant decrease in the surface quality of the steel sheet and adecrease in the ductility and weldability of the steel sheet.

Therefore, in the present disclosure, the content of Si may preferablybe within the range of 0.01 to 1.0%. More preferably, the content of Simay be within the range of 0.05% or more.

Manganese (Mn): 1.5% to 2.5%

Like Si, manganese (Mn) is an element effective in the solid-solutionstrengthening of steel, and increases the hardenability of steel tofacilitate the formation of bainite during cooling after heat treatment.

If the content of Mn is lower than 1.5%, the above-mentioned effectsobtainable by the addition of Mn are insufficient. Conversely, if thecontent of Mn exceeds 2.5%, martensitic phase transformation occurseasily due to a significant increase in hardenability, segregationregions are greatly developed in a thicknesswise center region of a slabduring slab casting in a continuous casting process, and impactresistance decreases due to the formation of thicknesswise non-uniformmicrostructures during cooling after hot rolling.

Therefore, in the present disclosure, the content of Mn may preferablybe within the range of 1.5% to 2.5%. More preferably, the content of Mnmay be within the range of 1.6% to 2.1%.

Aluminum (Al): 0.01% to 0.1%

Aluminum (Al) refers to sol.Al, and Al is an element mainly added fordeoxidation. If the content of Al is lower than 0.01%, the effect of Alis insufficient, and if the content of Al exceeds 0.1%, Al combines withnitrogen to form an AlN precipitate, facilitating the formation ofcorner cracks in a slab during continuous casting and the formation ofinclusion-induced defects.

Therefore, in the present disclosure, the content of Al may preferablybe within the range of 0.01% to 0.1%.

Chromium (Cr): 0.005% to 1.0%

Chromium (Cr) induces the solid-solution strengthening of steel and hasa function of helping bainite transformation at a coiling temperature bydelaying ferrite transformation during cooling.

If the content of Cr is lower than 0.005%, the effect of Cr may not besufficient. Conversely, if the content of Cr exceeds 1.0%, ferritetransformation is excessively delayed, and thus martensite is formed,resulting in poor elongation. In addition, similar to the case of Mn,segregation regions are greatly developed in a thicknesswise centerregion, and impact resistance decreases due to thicknesswise non-uniformmicrostructures.

Therefore, in the present disclosure, the content of Cr may preferablybe within the range of 0.005% to 1.0%. More preferably, the content ofCr may be within the range of 0.3% to 0.9%.

Phosphorus (P): 0.001% to 0.05%

Like Si, phosphorus (P) is an element having a solid-solutionstrengthening effect and a ferrite transformation promoting effect.

Maintaining the content of P within the range of less than 0.001% iseconomically disadvantageous due to excessive manufacturing costs andalso makes it difficult to guarantee strength. Conversely, if thecontent of P exceeds 0.05%, embrittlement may occur due to grainboundary segregation, fine cracks are easily formed during a formingprocess, and ductility and impact resistance are greatly deteriorated.

Therefore, in the present disclosure, the content of P may preferably bewithin the range of 0.001% to 0.05%.

Sulfur (S): 0.001% to 0.01%

Sulfur (S) is an impurity present in steel, and if the content of Sexceeds 0.01%, S combines with Mn or the like to form non-metallicinclusions, thereby facilitating the formation of fine cracks during asteel cutting process and markedly decreasing impact resistance.Conversely, to maintain the content of S within the range of less than0.001%, an excessive amount of time is required for a steel makingprocess, causing a decrease in productivity.

Therefore, in the present disclosure, the content of S may preferably bewithin the range of 0.001% to 0.01%.

Nitrogen (N): 0.001% to 0.01%

Together with C, nitrogen (N) is representative of solid-solutionstrengthening elements, and forms coarse precipitates by combining withTi or Al. Although the solid-solution strengthening effect of N issuperior to the solid-solution strengthening effect of carbon, thetoughness of steel markedly decreases as the content of N in the steelincreases, and thus it is preferable to adjust the content of S to 0.01%or less. To maintain the content of N within the range of less than0.001%, an excessive amount of time is required for a steel makingprocess, and thus productivity decreases.

Therefore, in the present disclosure, the content of N may preferably bewithin the range of 0.001% to 0.01%.

Niobium (Nb): 0.005% to 0.06%

Niobium (Nb), representative of precipitation enhancing elements, iseffective in improving the strength and impact toughness of steelbecause Nb precipitated during hot rolling delays recrystallization andthus causes grain refinement.

If the content of Nb is lower than 0.005%, the above-described effectsmay not be sufficient, whereas if the content of Nb exceeds 0.06%,formability and impact resistance are deteriorated by the formation ofelongated grains and coarse composite precipitates caused by anexcessive delay of recrystallization during hot rolling.

Therefore, in the present disclosure, the content of Nb may preferablybe within the range of 0.005% to 0.06%. More preferably, the content ofNb may be within the range of 0.01% to 0.05%.

Titanium (Ti): 0.005% to 0.11%

Together with Nb, titanium (Ti) is representative of precipitationenhancing elements and forms coarse TiN in steel due to strong affinitywith nitrogen (N). TiN has an effect of suppressing the growth of grainsduring heating for hot rolling. In addition, Ti remaining after reactionwith nitrogen dissolves in steel and bonds with carbon (C) to form a TiCprecipitate, and thus Ti is useful for improving the strength of steel.

If the content of Ti is lower than 0.005%, the above-described effectsmay not be sufficient, whereas if the content of Ti exceeds 0.11%,impact resistance for a forming process may be deteriorated due to theformation of coarse TiN and the coarsening of precipitates.

Therefore, in the present disclosure, the content of Ti may preferablybe within the range of 0.005% to 0.11%. More preferably, the content ofTi may be within range of 0.01% to 0.1%.

The remaining component of the steel sheet of the present disclosure isiron (Fe). However, impurities of raw materials or manufacturingenvironments may be inevitably included in the high-strength steelsheet, and such impurities may not be removed from the high-strengthsteel sheet. Such impurities are well-known to those of ordinary skillin manufacturing industries, and thus specific descriptions of theimpurities will not be given in the present specification.

It is preferable that the steel sheet of the present disclosuresatisfying the above-described alloy composition include a complex phaseof ferrite and bainite as a matrix thereof.

In this case, it is preferable that the sum of the area fractions offerrite and bainite be 90% or more. If the sum of the area fractions offerrite and bainite is lower than 90%, high strength and impactresistance may not be guaranteed.

In the composite phase, ferrite may preferably be included in an areafraction of 10% to 80%.

In addition, it is preferable that the ferrite phase have an averagegrain diameter (equivalent circular diameter) of 1 μm to 5 μm. If theaverage grain diameter of the ferrite phase exceeds 5 μm, there is aproblem in that formability and collision properties decrease due to theformation of non-uniform microstructures. Conversely, if the averagegrain diameter is maintained within the range of less than 1 μm, a lowrolling temperature or a high rolling reduction ratio is required,thereby causing an increase in the rolling load during a manufacturingprocess, excessive formation of shear texture, and significantdeterioration in formability.

Phases other than the complex phase may include a MA phase (a mixture ofmartensite and austenite) and a martensite phase, and it is preferablethat the total area fraction of the MA phase and the martensite phase be1% to 10%. If the total area fraction of the MA phase and the martensitephase is greater than 10%, tensile strength increases, but yieldstrength decreases, thereby causing cracks along the interface betweenthe MA phase and the martensite phase during collision and accordinglydecreasing impact resistance characteristics.

As described above, the total area fraction of the coarse MA phase andthe martensite phase in the matrix of the high-strength steel sheet isminimized such that unevenness in microstructure may be reduced.

Particularly, in the steel sheet of the present disclosure, the sheartexture ({110}<112>, {112}<111>) area ratio of a center region and asurface region may preferably be within the range of 0.05 to 1.0, wherethe surface region refers to a region from a surface of the steel sheetto 1/10t in a thickness direction of the steel sheet (here, “t” refersto the thickness (mm) of the steel sheet), and the center region refersto a region from 1/10t to ½t in a thickness direction of the steelsheet.

The texture {110}<112> and {112}<111> is mainly formed by the phasetransformation of a microstructure, non-recrystallized during hotrolling, and may be formed due to large shear deformation when thefriction between rolling rolls and a raw material (steel) in surfaceregions increases due to a low rolling temperature or a high rollingreduction. Such textures increase the anisotropy of the steel sheet andthus deteriorate the formability of the steel sheet. However, if theamount of a microstructure non-recrystallized during hot rollingincreases, ferrite deformation during phase tarnsformation isfacilitated, and thus a fine and uniform microstructure may be obtained.

Therefore, if the texture area ratio of the center region and thesurface region of the steel sheet is adjusted to be 0.05 to 1.0, auniform and fine microstructure is formed, and the fraction of ferriteincreases to relatively reduce the fractions of MA phase and martensitesuch that the yield strength and strain-rate sensitivity of the steelsheet may increase to increase energy absorption during a high-speedcollision. If the area ratio of the texture exceeds 1.0, theabove-described effects are saturated, and microstructures elongated inthe rolling direction increase significantly to cause an increase in theanisotropy of the steel sheet and a decrease in the formability of thesteel sheet.

Although a method of measuring the area fraction of the texture is notparticularly limited, for example, the area fraction of the texture maybe analyzed using Electron Back Scattered Diffraction (EBSD).Specifically, the area fractions of crystal orientations (110)[1 −1 −2]and (112) [−1 −1 1] in the center region and the surface region may beobtained from EBSD analysis on a rolled section, and the ratio of thearea fractions may be calculated.

The steel sheet of the present disclosure having the above-describedalloy composition and microstructure may have a tensile strength of 780MPa or more and may absorb energy at a rate of 80 J/m³ or more during acollision, that is, may have high strength and high impact resistance.

Hereinafter, a method for manufacturing a high-strength hot-rolled steelsheet having high bendability and high low-temperature toughness will bedescribed in detail according to another aspect of the presentdisclosure.

The high-strength steel sheet of the present disclosure may bemanufactured by performing a series of process [reheating-hotrolling-cooling-coiling] on a steel slab satisfying the alloycomposition proposed in the present disclosure.

Hereinafter, conditions of each of the processes will be described indetail.

[Reheating Steel Slab]

In the present disclosure, prior to hot rolling, it is preferable toreheat the steel slab so as to homogenize the steel slab, and in thiscase, it is preferable to perform the reheating process at 1200° C. to1350° C.

If the reheating temperature is lower than 1200° C., precipitates arenot sufficiently redissolved to result in a decrease in the formation ofprecipitates in the processes after the hot rolling process, and thereis a problem in that coarse TiN remains. Conversely, if the reheatingtemperature exceeds 1350° C., strength undesirably decreases due toabnormal growth of austenite grains.

[Hot Rolling]

Preferably, the reheated steel slab is hot rolled to produce ahot-rolled steel sheet, and in this case, the hot rolling process isperformed within the temperature range of 800° C. to 1150° C. underfinish hot rolling conditions satisfying [Equation 1] and [Equation 2].

When the hot rolling process starts at a temperature higher than 1150°C., the temperature of the hot-rolled steel sheet is increased, causingthe formation of coarse grains and deteriorating the surface quality ofthe hot-rolled steel sheet. In addition, if the hot rolling process endsat a temperature lower than 800° C., elongated crystal grains developdue to an excessive recrystallization delay, resulting in severeanisotropy and poor formability.

In particular, in the hot rolling process of the present disclosure, ifrolling is terminated at a temperature higher than the temperature rangesuggested in [Equation 1] below (that is, a temperature higher than Tn),the microstructure of the steel sheet is coarse and non-uniform, andphase transformation is delayed, thereby deteriorating impact resistancedue to the formation of coarse MA phase and martensite. In addition, ifrolling is terminated at a temperature lower than the temperature rangesuggested in [Equation 1] below (that is, a temperature less thanTn-50), the microstructure of the steel sheet is deformed due to anexcessive delay of recrystallization, and thus the formability of thesteel sheet may markedly deteriorate due to the formation of excessiveshear texture {110}<112> and {112}<111>.

The above-described microstructure changes may occur to guaranteeformability and impact resistance when [Equation 2] below is satisfied.More specifically, if the reduction of hot rolling is greater than therange suggested in [Equation 2] below (greater than 1.5×Ec), sheartexture is excessively formed, and formability deteriorates. Conversely,if the reduction of hot rolling is lower than the range suggested in[Equation 2] below (less than Ec), fine and uniform ferrite may not beformed during cooling immediately after hot rolling, and non-uniform andcoarse MA phase and martensite are formed because it is difficult tofacilitate ferrite transformation, thereby deteriorating impactresistance.

Tn-50≤FDT≤Tn   [Equation 1]

Where Tn refers to a temperature at which recrystallization delaystarts, Tn=730+92×[C]+70×[Mn]+45×[Cr]+780×[Nb]+520×[Ti]−80×[Si] whereeach element refers to a weight content, and FDT refers to thetemperature (° C.) of the hot-rolled steel sheet immediately afterfinish hot rolling.

Ec≤E1+E2≤1.5×Ec   [Equation 2]

Where Ec=4.75×10⁻⁴ (125×Exp(Qdef/(Rx(273+FDT))))^(0.17),Qdef=277000−2535×[C]+1510×[Mn]+9621×[Si]+1255×[Cr]+53680×[Ti]^(0.592)+70730×[Nb]^(0.565)where each element refers to a weight content.

In addition, E1 refers to the rolling reduction at the final pass of hotrolling, E2 refers to the rolling reduction before the final hot rollingpass (the pass just before the final pass), and R refers to the gasconstant: 8.314.

In the present disclosure, when the proposed alloy composition andmanufacturing conditions, particularly, hot-rolling conditions satisfy[Equation 1] and [Equation 2] at the same time, the strength,formability, and impact resistance of the steel sheet may be guaranteedas intended. More specifically, optimizing the alloy composition andmanufacturing conditions (hot rolling conditions) as described aboveleads to an effective delay of recrystallization during hot rolling,thereby promoting ferrite transformation during phase transformation,forming fine and uniform grains, improving strength and impactresistance.

In addition, owing to the promotion of ferrite transformation, theamounts of untransformed phases are reduced during the subsequentcooling process, thereby reducing the fractions of coarse MA phase andmartensite and preventing the formation of non-uniform microstructures.

If recrystallization is excessively delayed, there is a problem in thatdeformed microstructures are densely present in the entirety of therolled steel sheet or only in surface regions of the rolled steel sheetto cause deterioration of formability. However, according to the presentdisclosure, finish hot rolling is completed between therecrystallization delay start temperature Tn and Tn-50, and thus theabove-described effects may be obtained.

In addition, to effectively delay recrystallization and form uniform andfine ferrite grains, it may be necessary to adjust the reduction of hotrolling to be greater than a specific value while maintaining the hotrolling temperature within the above-described temperature range, buteven in this case, excessive reduction leads to a decrease informability. Therefore, it is necessary to adjust the reduction of hotrolling according to [Equation 2] proposed in the present disclosure.

[Cooling and Coiling]

The hot-rolled steel sheet manufactured as described above is cooledpreferably at an average cooling rate of 10° C./s to 100° C./s to atemperature range of 400° C. to 500° C. and is coiled with thetemperature range. In this case, the temperature range is based on thetemperature of the hot-rolled steel sheet.

At this time, if the cooling end temperature (coiling temperature)exceeds 500° C., pearlite is formed which makes it difficult toguarantee a target level of strength, and if the cooling end temperatureis lower than 400° C., martensite is excessively formed whichdeteriorates formability and impact resistance.

In addition, if the average cooling rate is lower than 10° C./s duringthe cooling within the above-mentioned temperature range, there is aproblem in that the grains of the matrix of the steel sheet is coarseand the microstructure of the steel sheet is uneven, and if the averagecooling rate exceeds 100° C./s, a MA phase is easily formed whichdeteriorates formability and impact resistance.

[Final Cooling]

A coil obtained by cooling and coiling the steel sheet as describedabove may preferably be cooled to the temperature range of roomtemperature to 200° C. at a cooling rate of 0.1° C./hour to 25° C./hour.

In this case, if the cooling rate exceeds 25° C./hour, someuntransformed phases in the steel sheet may easily be transformed into aMA phase, thereby deteriorating formability and impact resistance. Inaddition, adjusting the cooling rate to be less than 0.1° C./hour isuneconomical because additional heating equipment is required.

After this final cooling, the steel sheet is pickled and oiled, and thenheated to a temperature range of 450° C. to 740° C. to perform a hot-dipgalvanizing process.

The hot-dip galvanizing process may use a zinc-based plating bath, andalthough the alloy composition in the zinc-based plating bath is notparticularly limited, for example, the zinc-based plating bath may havean alloy composition of magnesium (Mg): 0.01% to 30% by weight, aluminum(Al): 0.01% to 50% by weight, and the balance of Zn and inevitableimpurities.

Hereinafter, the present disclosure will be described in more detailthrough an example. However, it should be noted that the followingexample is for illustrative purposes only and is not intended to limitthe scope of the present disclosure. The scope of the present disclosureis defined by the appended claims, and modifications and variationsreasonably made therein.

Mode for Invention EXAMPLE

Steel slabs having the alloy compositions shown in Table 1 below wereprepared. Here, the alloy compositions are given in weight %, and theremainder include Fe and inevitable impurities. Steel sheets weremanufactured from the prepared steel slabs under the manufacturingconditions shown in Table 2 below. In Table 2 below, FDT denotes atemperature during finish hot rolling, and CT denotes a coilingtemperature.

In addition, the reheating temperature of the steel slabs was 1200° C.,the thickness of hot-rolled steel sheets after hot rolling was 3mm,cooling immediately after hot rolling was performed at a cooling rate of20° C./s to 30° C./s, and the cooling rate after coiling was constant at10° C./hour.

TABLE 1 Steels C Si Mn Cr Al P S N Ti Nb CS1 0.15 0.03 1.7 0.4 0.03 0.010.003 0.004 0.08 0.03 CS2 0.08 0.4 1.8 0.8 0.03 0.01 0.003 0.004 0.090.02 CS3 0.07 0.5 2.1 0.5 0.04 0.01 0.002 0.005 0.10 0.04 CS4 0.08 0.11.8 0.01 0.03 0.01 0.003 0.004 0.09 0.04 CS5 0.11 0.2 2.0 0.3 0.03 0.010.003 0.004 0.08 0.03 CS6  0.095 0.3 1.9 0.4 0.03 0.01 0.003 0.003 0.030.05 CS7 0.12 0.5 1.9 0.7 0.04 0.01 0.003 0.003 0.10 0.02 CS8 0.13 0.11.8 0.8 0.04 0.01 0.003 0.003 0.07 0.05 IS1  0.065 0.3 1.8 0.5 0.03 0.010.003 0.004 0.08 0.03 IS2 0.07 0.01 1.6 0.7 0.03 0.01 0.003 0.0042 0.090.03 IS3 0.07 0.9 1.7 0.6 0.03 0.01 0.003 0.0035 0.07 0.035 IS4 0.07 0.31.6 0.7 0.03 0.01 0.003 0.004 0.10 0.02 IS5 0.10 0.7 2.0 0.9 0.03 0.010.003 0.004 0.03 0.025 IS6 0.11 0.06 2.0 0.9 0.03 0.01 0.003 0.004 0.040.04 IS7 0.13 0.6 1.95 0.6 0.03 0.01 0.003 0.003 0.05 0.03

TABLE 2 Equation 1 Equation 2 FDT CT Evalu- Evalu- Steels (° C.) (° C.)Tn ation E1 E2 Ec Qdef ation CS1 902 446 943 ∘ 9 20 20.6 301766 ∘ CS2940 435 930 x 9 18 18.5 305028 ∘ CS3 874 455 949 x 13 23 27.4 310641 ∘CS4 896 490 934 ∘ 15 25 22.3 304869 x CS5 906 485 943 ∘ 6 13 21.0 303831x CS6 850 453 920 x 16 26 26.8 302768 x CS7 892 530 933 ∘ 9 25 23.5306745 ∘ CS8 930 385 971 ∘ 8 25 19.4 305492 x IS1 876 445 925 ∘ 10 2124.5 304856 ∘ IS2 911 459 949 ∘ 10 18 20.2 302871 ∘ IS3 850 452 874 ∘ 1527 30.8 310563 ∘ IS4 880 455 924 ∘ 12 21 23.9 304495 ∘ IS5 879 445 899 ∘10 23 23.4 303164 ∘ IS6 922 430 968 ∘ 8 18 18.6 300908 ∘ IS7 885 457 907∘ 12 21 23.6 305006 ∘(In Table 2 above, Ec refers to calculated values expressed inpercentages (calculated value×100).)

The mechanical properties, that is, tensile strength (TS), yieldstrength (YS), and elongation (T-E1), and the absorbed impact energy ofeach of the steel sheets prepared as described above were measured, andthe microstructure of the steel sheet was observed as shown in Table 3below.

Specifically, the yield strength and elongation respectively refer toyield strength at 0.2% offset and breaking elongation, and the tensilestrength was measured using a JIS No. 5 specimen taken in a directionperpendicular to the rolling direction.

In addition, the absorbed impact energy was measured by a high-speedtensile test method, and since the strain rate of a material is from 100s⁻¹ to 500 s⁻¹ when an automobile crashes, a high-speed tensile test wasperformed at strain rates of 200 s⁻¹ and 500 s⁻¹ on a specimen takenaccording to the same standard as the above-mentioned tensile specimento measure the area under the obtained stress-strain curve until anelongate of 10% as the absorbed impact energy of the specimen.

In addition, the MA phase formed in the steel sheet was analyzed byetching the steel sheet using a Lepera etching method and observing thesteel sheet at a magnification of 1000 times using an optical microscopeand an image analyzer, and austenite was analyzed using Electron BackScattered Diffraction (EBSD) at a magnification of 3000 times. Inaddition, the fractions of martensite, ferrite, and bainite wereanalyzed using a scanning electron microscope (SEM) at magnifications of3000 times and 5000 times. In addition, the area fraction of sheartexture was measured using the above-mentioned Electron Back ScatteredDiffraction (EBSD).

TABLE 3 Microstructure Mechanical Shear Impact properties Textureresistance YS TS T − El F B γ M Area 200⁻¹ 500⁻¹ Steels (MPa) (MPa) (%)(%) (%) (%) (%) fraction (J/m³) (J/m³) CS1 725 978 9 45 37 4 14 0.85 7781 CS2 688 855 14 61 28 3 8 1.05 78 86 CS3 728 924 9 58 34 1 7 1.40 7483 CS4 695 870 10 59 34 0 7 1.35 77 88 CS5 740 956 10 33 56 2 9 0.31 7887 CS6 755 960 8 48 44 2 6 1.50 76 79 CS7 688 815 15 57 32 2 1 0.36 7283 CS8 885 1085 7 22 53 4 21 0.14 78 80 IS1 685 815 14 74 23 1 2 0.35 8293 IS2 689 792 15 72 26 0 2 0.28 81 94 IS3 675 821 15 76 21 1 2 0.33 8292 IS4 715 835 17 75 21 1 3 0.34 83 90 IS5 803 922 11 38 56 2 4 0.30 95104 IS6 820 994 10 16 76 2 6 0.11 96 102 IS7 815 1008 10 14 79 2 5 0.1394 110(In Table 3, γ and M respectively refer to fractions in the MA phase.)

As shown in Tables 1 to 3 above, each of Inventive Steels 1 to 7satisfying both the alloy composition and manufacturing conditionsproposed in the present disclosure had a complex phase of ferrite andbainite as a matrix, and a shear texture area ratio, that is, a sheartexture area ratio of a center region and a surface region (sheartexture area fraction in center region/shear texture area fraction insurface region) within the range of 0.05 to 1.0. Thus, Inventive Steels1 to 7 had intended high strength and large amounts of absorbed impactenergy.

However, in Comparative Steel 1 having an excessive C content comparedwith the alloy composition proposed in the present disclosure, MA phaseand martensite were excessively formed because of a high C content in anuntransformed phase, and thus absorbed impact energy of ComparativeSteel 1 was low. It is considered that the reason for this is rapidfracture along the boundary between MA phase and martensite duringhigh-speed deformation.

Comparative steels 2 and 3 are cases in which the finish hot rollingtemperature during hot rolling does not satisfy Equation 1. Particularlyin Comparative Steel 2 of which FDT exceeds the range specified inEquation 1, almost no shear texture was developed, and ferritetransformation was not promoted, resulting in the formation of MA phaseand martensite in high fractions and a small amount of absorbed impactenergy. Furthermore, in Comparative Steel 3 of which FDT is excessivelyless than the range specified in Equation 1, the amount of shear texturein the surface region was more increased than the center region, andthus absorbed impact energy was low due to a significant decrease inelongation even though strength was increased.

Comparative steels 4 and 5 are cases in which the rolling reductionconditions for hot rolling do not satisfy Equation 2. In ComparativeSteel 4 of which the sum of the reduction in the final pass and thereduction in the previous pass exceeds the range specified in Equation2, the amount of shear texture in the surface region was more increasedthan the center region, and thus absorbed impact energy was low due to asignificant decrease in elongation even though strength was increased.Furthermore, in Comparative Steel 5 of which the sum of the reduction inthe final pass and the reduction in the previous pass is lower than therange specified in Equation 2, ferrite transformation was not promoted,and the fractions of MA phase and martensite phase were markedlyincreased, resulting in a small amount of absorbed impact energy.

In Comparative Steel 6 of which the hot rolling conditions do notsatisfy both of Equations 1 and 2, the amount of shear texture in thesurface region was excessively greater than the amount of shear texturein the center region, resulting in a significant decrease in elongationand a decrease in absorbed impact energy. In particular, the absorbedimpact energy at a condition of 500 s⁻¹ was markedly decreased, and thusComparative Steel 6 was not suitable for collision at a high strainrate.

Comparative steels 7 and 8 are cases in which the temperature rangeduring coiling is outside the range proposed in the present disclosure.In Comparative Steel 7 having an excessively high coiling temperature,pearlite was formed in the microstructure, resulting in failure inguaranteeing a target level of strength and a small amount of absorbedimpact energy. It is considered that the reason for this is that: sincenot bainite but perlite was formed from an untransformed phase due to ahigh coiling temperature, the strength of steel was not sufficientlysecured, and fracture rapidly propagated to the pearlite duringhigh-speed deformation. Comparative steel 8, of which the coilingtemperature was low, had excessively high strength due to the formationof a large amount of martensite and absorbed a small amount of impactenergy. It is considered that the reason for this is that duringhigh-speed deformation, fracture rapidly occurred because of a localincrease in dislocation density.

FIG. 1 is images of shear texture in a surface region (A) and sheartexture in a center region (B) of Inventive Steel 6, wherein the areafraction of the shear texture in the surface region (A) is0.217(0.130+0.087), the area fraction of the shear texture in the centerregion (B) is 0.025(0.019+0.006), and the ratio (B/A) is 0.11.

1. A method for manufacturing a steel sheet, the method comprising:reheating a steel slab at a temperature of 1200° C. to 1350° C., thesteel slab comprising, by weight%, carbon (C): 0.05% to 0.14%, silicon(Si): 0.01% to 1.0%, manganese (Mn): 1.5% to 2.5%, aluminum (Al) : 0.01%to 0.1%, chromium (Cr): 0.005% to 1.0%, phosphorus (P): 0.001% to 0.05%,sulfur (S): 0.001% to 0.01%, nitrogen (N) : 0.001% to 0.01%, niobium(Nb): 0.005% to 0.06%, titanium (Ti): 0.005% to 0.11%, and a balance ofiron (Fe) and inevitable impurities; finish hot rolling the reheatedsteel slab under conditions satisfying [Equation 1] and [Equation 2]below to obtain a hot-rolled steel sheet; cooling the hot-rolled steelsheet at a cooling rate of 10° C./s to 100° C./s to a temperature of400° C. to 500° C. after the finish hot rolling; and coiling the steelsheet at a temperature of 400° C. to 500° C. after the cooling,Tn-50≤FDT≤Tn   [Equation 1] in which Tn refers to a temperature at whichrecrystallization delay starts,Tn=730+92×[C]+70×[Mn]+45×[Cr]+780×[Nb]+520×[Ti]−80×[Si] where eachelement refers to a weight content, and FDT refers to a temperature (°C.) of the hot-rolled steel sheet immediately after the finish hotrolling,Ec≤E1+E2≤1.5×Ec   [Equation 2] in which Ec=4.75×10⁻⁴(125×Exp(Qdef/(R×(273+FDT))))^(0.17),Qdef=277000−2535×[C]+1510×[Mn]+9621×[Si]+1255×[Cr]+53680×[Ti]^(0.592)+70730×[Nb]^(0.565)where each element refers to a weight content, and E1 refers to arolling reduction in a final pass of hot rolling, E2 refers to a rollingreduction before the final hot rolling pass (pass just before the finalpass), and R refers to the gas constant: 8.314.
 2. The method of claim1, further comprising cooling the coiled steel sheet at a rate of 0.1°C./hour to 25° C./hour to a temperature range of room temperature to200° C.
 3. The method of claim 2, further comprising pickling and oilingthe coiled steel sheet after the cooling.
 4. The method of claim 3,further comprising heating the steel sheet to a temperature range of450° C. to 740° C. after the pickling and oiling, and then hot-dipgalvanizing the steel sheet.
 5. The method of claim 4, wherein thehot-dip galvanizing is performed using a plating bath comprising, byweight%, magnesium (Mg): 0.01% to 30%, aluminum (Al): 0.01% to 50%, andthe balance zinc (Zn) and inevitable impurities.
 6. The method of claim1, wherein the steel sheet has a microstructure comprising ferrite andbainite in a total area fraction of 90% or more, a MA phase, which is acomplex phase of martensite and austenite, and martensite in a totalarea fraction of 1% to 10%, and the steel sheet has a value of 0.05 to1.0 as a {110}<112>, and {112}<111>—shear texture area ratio of a centerregion ranging deeper than 1/10t to ½t in a thickness direction, trefers to thickness (mm) and a surface region ranging from a surface to1/10t in the thickness direction.
 7. The method of claim 6, wherein theferrite has an average grain diameter within a range of 1 μm to 5 μm. 8.The method of claim 6, wherein the steel sheet has a tensile strength of780 MPa or more, and absorbs energy in an amount of 80 J/m3 or moreduring a collision.